KR920001390B1 - Surface-coated cemented carbide and a process therefor - Google Patents

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Description

Coated cemented carbide and its manufacturing method

1 is a graph showing the surface hardness distribution of alloy samples 1 to 4 according to a preferred embodiment of the present invention.

2 is a graph showing the distribution of CO on the surfaces of alloy samples 8 to 11 according to a preferred embodiment of the present invention.

The present invention relates to a coated cemented carbide alloy having a very high toughness used in cutting tools of the present invention. It is about.

Recently, N / C machines have been introduced into the field of cutting, and factory automation (FA) has been advanced significantly. In this case, since the reliability of the cutting tool becomes very important, the development of a cutting tool having a higher toughness than the conventional one is required. In order to satisfy this requirement, cemented carbide (WP 159299/1977 and 194239/1982) consisting of WC-CO only with a surface layer, and a method of enriching the alloy surface with CO (Japanese special) Points 105628/1987, 187678/1985 and 194239/1982, i.e. U.S. Patent Nos. 4, 610, 931) and a method for the presence of free carbon in the alloy to prevent the formation of a decarburized layer underneath the coating layer. 155190/1977) and the like have been proposed.

However, cemented carbides with a WC-CO layer only on the surface or a CO-enriched layer on the surface have improved toughness but are problematic for wear resistance. Particularly, under conditions of high cutting speed, an alloy having an alloy enriched with CO may not be able to withstand actual use due to the fast wear rate of the rake face. In the case of an alloy containing free carbon [FC], the toughness is improved by increasing the amount of carbon, but when it exceeds 0.2% by weight, the alloy is agglomerated to lower the strength of the alloy itself.

It is an object of the present invention to provide a novel coated cemented carbide to solve the above problems of the prior art.

Another object of the present invention is to provide a cemented carbide substrate having high toughness and high wear resistance.

Another object of the present invention is to provide a highly efficient cutting tool composed of a coated cemented carbide coated with a hard thin film such as titanium carbide.

Another object of the present invention is to provide a method for producing a coated cemented carbide.

These objectives are based on at least one hard phase selected from the group consisting of carbides, nitrides and carbonitrides of the Group IVa, Va, and VIa metals of the periodic table and the IVa of the periodic tables on a cemented carbide substrate composed of at least one bond phase selected from the iron group metal It is achieved by at least one selected from the group consisting of carbides, nitrides, oxides and borides of Group Va, Group VIa metals and their cemented carbides surface-coated with a single layer or multiple layers of a solid solution, aluminum oxide, wherein the coating layer and the substrate The hardness of the cemented carbide substrate is in the range of 2 to 5 µm from the interface between them, and the hardness of the cemented carbide substrate is 800 to 1300 kg / mm2 with the beaker hardness of 500 g load, and increases uniformly toward the inside of the substrate, and is constant within the range of approximately 500 to 100 µm from the interface. do.

Hereinafter, the present invention will be described in detail with reference to the accompanying drawings.

The present inventors have made various efforts to develop the surface coated cemented carbide material for cutting tools having the highest properties, i.e., the wear resistance by the coating layer, and having higher toughness than the alloy of the prior art. I found that.

(I) A substrate comprising at least one hard phase selected from the group consisting of carbides, nitrides and carbonitrides of the Group IVa, Va, and VIa metals of the periodic table, and at least one bond phase selected from the iron group metals. Or preferably consisting of WC and CO, or a mixed carbide or mixed carbonitride of W, Ti, Nb and / or Ta and CO, or more preferably 10 to 96% by weight of WC, 1 to A cemented carbide composed of 70% by weight of Ti, W, Ta and / or Nb mixed carbonitrides and 3 to 20% by weight of CO is used.

(II) In the vicinity of the surface of a cemented carbide substrate composed mainly of WC and CO and having a thickness of 5 to 10 탆, the range of 2 to 20 탆 from the interface, preferably 5 to 5 to 5 to 100 탆. The amount of the bonded phase of the cemented carbide substrate in the range is 1.5 to 7 times the weight of the average amount of the bonded phase present in the alloy, and in particular, the cemented carbide substrate within the range of 2 to 20 µm, preferably 2 to 10 µm, immediately below the interface. The amount of binding phase CO is 1.5 to 7 weight times for the range of approximately 50 to 100 mu m. The amount of bonding phase in the range from the interface to 5 mu m is smaller than the inside of the cemented carbide substrate, and more preferably the CO content of the cemented carbide substrate in the range from the interface to 3 mu m is smaller than the range below 3 mu m from the interface.

(III) The layer composed mainly of WC and CO near the surface of the cemented carbide substrate, in particular, the hardness within the range of 2 to 5 μm from the interface is 700 to 1300 kg / mm 2, preferably 800 to 800 Vickers (Hv) hardness under a 500 g load. 1300 kg / mm 2, more preferably 950-1250 kg / mm 2, most preferably 1000-1200 kg / mm 2, and the hardness of the substrate is constantly increasing toward its interior and is approximately 50-100 μm from the interface. It becomes constant, Preferably it is 1500-1700 kg / mm <2> by the beaker hardness under 500g load.

(IV) When the bonding phase is CO, the amount of free carbon [FC] in the cemented carbide is (FC / CO + FC) of 1 to 2.4% by weight based on the amount of CO. When the bonding phase is Ni, the amount of [FC] is based on Ni. As a result, (FC / Ni + FC) is 0.5 to 2.2% by weight.

(Ⅴ) The amount of free carbon [FC] and nitrogen [N] in the cemented carbide substrate has the following relationship.

In formula, [FC] and [N] are represented by weight%.

The coated cemented carbide of the present invention having the above-described structure and features, the starting material described in formula (I) is cooled at a cooling rate of 0.1 to 10 ℃ / min, preferably from 0.1 to 10 within the temperature range of 1310 to 1225 ℃ Sintering including cooling at a cooling rate of 10 ° C./min or performing cooling at a time period of 10 minutes to 15 hours within a temperature range of 1310 to 1225 ° C., and the resulting substrates are then subjected to IVa, Va, It is prepared by coating at least one selected from the group consisting of carbides, nitrides, oxides and borides of group VIa elements, with a single layer or multiple layers of a solid solution and aluminum oxide thereof.

Preferably, the cemented carbide substrate obtained by the sintering process described above is subjected to chemical, mechanical or electrochemical processing to move CO or CO and C from the surface portion of the cemented carbide substrate.

The characteristics and structure of the surface-coated cemented carbide of the present invention and a method of manufacturing the same will be described in detail below.

The cemented carbide substrate of the present invention contains at least one hard phase selected from the group consisting of carbides, nitrides and carbonitrides of Group IVa, Va, and VIa metals, so that the nitrogen-containing hard phases are denitrified in part of the sintering process. And decomposition forms, for example, a layer of mainly WC and CO when in the hard phase WC. "Predominantly" generally means that the nitrogen-containing hard phase does not completely decompose and a small amount of nitrogen remains. In this case, preferably, [FC] and [N] in the cemented carbide should satisfy the following relationship.

In formula, [FC] and [N] are represented by weight%. For example, if the analytical amounts of [FC] and [N] in the alloy are 0.1% and 0.03%, respectively, 0.1 + 12/14 × 0.03 = 0.12. In this formula, [FC] denotes the amount of free carbon in the bond phase and [N] denotes the amount of nitrogen in the cemented carbide. When cemented carbide is produced by sintering, CO and C form a CO-C melt through eutectic reaction at an eutectic temperature of approximately 1309 ° C. However, in actual cemented carbides, C and W are characterized by dissolving in CO to use CO-WC melt through eutectic reaction, and the efficient use of such melt is carried out within the above-described range (hereinafter referred to as carbon equivalent). While nitrogen is expected to exhibit carbon-like behavior.

The cemented carbide having the above-described composition is cooled at a cooling rate of 0.1 to 10 DEG C / min, 1310 to 1225 DEG C, preferably at 1310 to 1255 DEG C, and preferably at a cooling rate of 1 to 5 DEG C / min. . 1255 ° C is a eutectic temperature at which CO, C and η phase (η phase means CO, W and C compounds) coexist and are due to a significant decrease in the carbon content in the alloy surface. Cooling of the cemented carbide is performed for 10 minutes to 15 hours in the temperature range of 1310 to 1225 ° C.

In the case where the bonding phase is CO or Ni, the amount of [FC] in the alloy is preferably within the range in which the liquid phase of the CO-C eutectic composition or the Ni-C eutectic composition is exhibited, thereby achieving the object of the present invention. That is, the amount of [FC] is 1 to 2.4% by weight relative to the amount of CO in the case of a CO bond phase, and 0.5 to 2.2% by weight relative to the amount of Ni in the case of a Ni bond phase. If it is higher than the upper limit, the compound of CO or Ni and C is precipitated as an initial crystal and should be avoided. If it is lower than the lower limit, the eutectic liquid phase does not appear. In this case, the object of the present invention cannot be achieved.

Since the hard phase containing nitride as described in (I) is subjected to denitrification to reduce the carbon equivalent of the alloy surface, the CO-W-C melt in the alloy is transferred to these surfaces. In other words, the concentration of CO-W-C melt is generated on the surface of the alloy through diffusion of the CO-W-C melt, and the alloy strength is increased after sintering. In particular, since the alloy surface is mainly composed of WC-CO, generally WC- (4.5 to 60% by weight) CO, the strength is greatly reduced, and the beaker hardness of 500 g load is 700 to 1000 kg / mm 2. When the carbon amount described in (V) is 0.06 or less, the diffusion of the CO-WC melt is too small to achieve the structure of the present invention. On the other hand, when the carbon amount is 0.17 or more, the compounds of CO and C are columnar on the alloy surface. Precipitates to crystals and becomes brittle. If the temperature exceeds the above-mentioned range, that is, 1310 ° C., the moving speed of the CO-WC melt becomes large so that it flows out of the alloy surface and no constant change in hardness is given, whereas if it is below 1225 ° C., the CO-WC melt The hardness change described above cannot be given because it is not formed. If the cooling rate exceeds 10 ℃ / min, the movement of the CO ℃ -WC melt is too small to give a hardness change, while if less than 0.1 ℃ / min the commercial scale productivity is to be avoided to avoid the cooling rate is 1 ~ A 5 ° C./min range is preferred.

In the process of sintering the alloy, the denitrification reaction in the alloy is preferably suppressed by introducing, for example, N ₂ , CH ₄ , H ₂ , Ar gas, etc. up to 1310 ° C., and high vacuum within the range of 1310-1225 ° C. Or decarbonization or sintering under an oxidizing atmosphere such as H ₂ , H ₂ + H ₂ O, CO ₂ , CO ₂ + CO, and the like.

The alloy surface layer, which is mainly composed of WC and CO, is formed through decomposition of the nitride-containing hard phase, but may be formed by nitriding and denitrifying the IVa, Va or VIa metal during the temperature increase.

In the present invention, the hardness of the alloy surface is generally in the range of 700 to 1000 kg / mm 2, and if it is 700 kg / mm 2 or less, the toughness is remarkably improved, but the wear resistance is lowered, which is a problem in practical use, whereas 1000 kg / If it is mm <2> or more, toughness improvement cannot be expected. Surface hardness can be controlled by the cooling rate and the degree of denitrification or decarbonization of the alloy surface. In order to maintain satisfactory wear resistance and toughness, that is, from the viewpoint of widely using the alloy for various purposes, the hardness of the surface layer in the range of 2 to 5 μm from the inside is 700 to 1300 kg / mm 2, preferably 950 to 1250 kg / mm 2 And more preferably 1000 to 1200 kg / mm 2, and adjust the internal hardness in the range of approximately 50-100 μm from the alloy surface to 1500 to 1700 kg / mm 2, sometimes outside of this range with regard to widespread use. Becomes Hardness is the hardness of beakers under 500g load, and it is a general ceramic depending on the weight. The hardness of the surface layer is somewhat higher at the load of 500g or more.

When the cemented carbide substrate of the present invention is sintered by the above-described process, the amount of the binder phase in the alloy in the range of 50 to 100 µm to 2 to 20 µm from the interface between the alloy surface and the coating layer is 7 to 1.5 weight times the average amount of the bond phase. . In particular, the amount of the bonding phase in the range of 50 mu m from the alloy surface is more than three times, much higher than the prior art described in Japanese Patent Laid-Open No. 199239/1982. According to the present invention, the bonding phase in the alloy surface becomes much richer.

In the present invention, since CO or CO and C are present on the alloy surface, the cutting tool has a somewhat larger crater depth at higher cutting speeds even in actual cutting using the surface-coated cemented carbide as the cutting tool. A problem occurs. In this case, the problem can be solved by reducing or eliminating the bonding phase from the interface between the coating layer and the surface of the alloy to 5 µm, preferably in the range of 1 to 5 µm, smaller than the average amount of the bonding phase present in the alloy. This is because toughness is greatly lowered. In the case of dissipating the bonded phase, the range is preferably at most 3 µm, and when it exceeds 3 µm, the toughness is greatly lowered. The reduction or disappearance of the binding phase is carried out by chemical treatment with acids such as nitric acid, hydrochloric acid, hydrofluoric acid, sulfuric acid, and the like by mechanical treatment or electrochemical treatment such as barrel treatment and brushing.

The coating layer used in the present invention comprises at least one selected from the group consisting of carbides, nitrides, oxides and borides of elements IVa, Va, and VIa of the periodic table, and a single layer or a multilayer consisting of a solid solution and aluminum oxide thereof by CVD method. It is formed by coating and has a thickness of 1-20㎛.

The coated cemented carbide of the present invention has excellent abrasion resistance by the coating layer and higher toughness than the alloy of the prior art, thus providing a tool that is much more reliable than the prior art tools.

The following examples illustrate the invention in detail, and unless otherwise indicated, percentages represent weight percentages.

Example 1

2.5% Ti (CN), 3.0% TaC, 6.0% CO and WC residues were mixed to bring [FC] + 12/14 × [N] (carbon equivalent) in the alloy as shown in Table 1, in the heating in a vacuum to 1400 ℃ 2torr and the N ₂ atmosphere at a cooling speed for 30 minutes and then held at a cooling rate of 10 ℃ / min cooling to 1310 ℃, and 3 ℃ / min in a vacuum (10 ^-3 torr) to 1200 ℃ Cool. The resultant cemented carbide was coated with an inner layer of 5 µm TiC and an outer layer of 1 µm Al ₂ O ₃ by a conventional CVD method, and then subjected to a cutting test under the following conditions (type: CNMG 120408; holder type: PCLNR2525-43). .

For comparison, the same M20 claddings and those coated with 5 μm TiC, 1 μm Al ₂ O ₃ are tested identically.

Table 1 shows the test results and the Hv hardness of the substrate under a 500g load within a range of 5 μm from the inner layer.

Cutting condition A (wear resistance test)

Cutting speed: 180m / min

Feed: 0.36㎜ / rev

Depth of cut: 2.0㎜

Workpiece: SCM435

Cutting time: 20 minutes

Cutting condition B (toughness test)

Cutting speed: 60m / min

Feed: 0.20 to 0.40 mm / rev

Workpiece: SCM435 (grooved into 10㎜ × 50㎜)

Cutting time: 30 seconds 8 times

TABLE 1

As a result of observation of the cross-sectional structure of the alloy surface of Samples 1 to 4, only the WC-CO layer was formed in the range of approximately 5 µm from the surface, and mixed carbonitrides of [Ti, Ta, W] [CN] exist in the range of 5 µm. It was found that FC precipitates inside the alloy. 1 shows the hardness distribution of the surface layers of Samples 1-4. The hardness in the 100-micrometer range below the alloy surface was 1500 kg / mm 2.

In the following Example 2, an alloy in which CO or CO and C were transferred in a 0.5 μm range was immersed in a 10% nitric acid solution at 20 ° C. for 10 minutes.

Example 2

In order to sinter the sample 3 of Example 1, WC powders with grain sizes of 4 μm and 2 μm were sintered and coated using the ratios of 1: 1 and 1: 2 and tested in a similar manner as in Example 1.

As a result, in the test A, the former had a side wear width of 0.18 mm, the latter showed 0.15 mm, and in the test B, the former had a breakage rate of 8% and the latter of 12%. In the former case, the hardness of the alloy surface was 1070 kg / mm 2, the latter was 1120 kg / mm 2, and in the range of 100 μm from the alloy surface, the former was 1600 kg / mm 2 and the latter was 1680 kg / mm 2.

Example 3

The fourth sintered compact of Example 1 was immersed in a 10% aqueous nitric acid solution for 10 minutes, 25 minutes in the same solution (ii), and 20 minutes in 20% nitric acid solution at a temperature maintained at 20 DEG C. Samples 5-7 are made by moving CO and C on the alloy surface.

These alloys were coated in a similar manner to Example 1, tested A and B, and the results are shown in Table 2.

TABLE 2

In the case of sample 5, the amount of CO in the range from the surface to 2 µm is smaller than the inside, and in the case of samples 6 and 7, CO is extinguished in the range from 5 to 3 µm from the surface, respectively.

Example 4

An alloy consisting of 2.0% Ti (CN), 3% TaC 5.6% CO and WC residue and having a carbon equivalent of 0.15 was sintered in a similar manner to Example 1 and cooled at 1310 ° C. and shown in FIG. It cooled to 1200 degreeC under the conditions.

TABLE 3

CO enrichment near the surface of the alloy was analyzed by EPMA (ACC: 20 kv, SC: 200 A, beam diameter: 10 μm) to obtain the results shown in FIG.

Example 5

An alloy consisting of 2.5% Ti (CN), 6.0% TaC, 5.6% CO and WC residue and having a carbon equivalent of 0.15 was heated to 1400 ° C. under vacuum and 2 ° C./min in a CH ₄ and H ₂ atmosphere. After cooling to 1310 ⁸ at a cooling rate, it is cooled to 1200 ° C. at a cooling rate of 0.5 ° C./min in a vacuum (10 ⁻⁵ torr) or CO ₂ atmosphere. The resulting alloy has a surface hardness of 920 kg / mm 2, and the hardness is constantly increased within the range of 70 µm below the surface to a constant value of 1600 kg / mm 2. Within the range of 5 mu m from the surface, the amount of mixed carbonitrides of [Ti, Ta, W] [CN] is reduced compared to the inside of the alloy.

When this alloy was coated with a layer of ₃ μm TiC, 2 μm TiN, 1 μm TiCN, and 1 μm of Al ₂ O ₃ and subjected to a cutting test in the same manner as in Example 1, the side wear width was 0.23 mm and the failure rate was 3 %to be.

Example 6

An alloy consisting of 2.0% Ti (CN), 6.0% TaC, 5.6% CO and WC residue and having a carbon equivalent of 0.15 was sintered in a similar manner as in Example 1 and cooled to 1310 ° C., followed by the conditions shown in Table 4. Cooled to 1200 ° C.

TABLE 4

Example 7

The sample sixteenth alloy of Example 6 was immersed in 1.0% aqueous nitric acid solution for 10 minutes, neutralized with 5% aqueous sodium hydroxide solution for 5 minutes, washed with water for 5 minutes, sprayed with diamond grain No. 1000, and cleaned with a steel brush. The alloy thus treated was coated with a 5 μm TiC, 1 μm Al ₂ O ₃ layer to perform a cutting test in a similar manner to Example 1. Samples that were not iron oxide initially peeled off, while acid treated samples exhibited normal wear.

Example 8

Alloy powder consisting of 2.0% TiC, 6.0% TaC, 5.6% CO and WC residue was formed in shape No.SNG432, heated to 1000 ° C. in vacuum, and sintered at 1000 to 1450 ° C. in N ₂ atmosphere to produce alloy carbon. The equivalent weight of 0.15 was then cooled in a similar manner to Example 5 to obtain an alloy having a structure and hardness distribution substantially similar to that of Example 5.

Example 9

Alloy powder consisting of 2.0% Ti (CN), 5.0% TaC, 5.6% CO and WC residue. Formed with SNG432, heated in vacuo and sintered at 1400 ° C. in vacuo to have a carbon equivalent of 0.15. The resultant alloy was processed to a predetermined shape and subjected to edge formation, heated to 1350 ° C., held at 5 torr in N ₂ atmosphere for 30 minutes, rapidly cooled to 1310 ° C. at a cooling rate of 20 ° C./min, and then cooled to 10 ⁻⁵ . Cooled to 1200 ° C. at 1310 ° C. at 2 ° C./min in vacuo.

The resulting alloy has a WC-CO layer in the range of 2 μm from the alloy surface and has a surface hardness of 1020 kg / mm 2. Similarly, when sintering at 0.5torr in a CO ₂ atmosphere, the surface hardness was 990kg / mm 2.

Example 10

The composition similar to Example 1 is mixed with the amount of free carbon based on the amount of CO as 1, 1.5, 2, 2.4%. When the resulting alloy was tested under cutting condition B, the failure rates were 23%, 8%, 2% and 0%, respectively.

Example 11

The sample fourth alloy of Example 1 was dipped in 20% aqueous nitric acid solution at 20 ° C. for 20 minutes, 10 minutes, and 5 minutes. In the sample treated for 20 minutes, the CO phase disappeared within the range of 5 µm from the surface, and in the case of the sample treated for 10 minutes, the CO phase disappeared within the range of 3 µm from the surface, and in the sample treated for 5 minutes, 1 The CO phase disappears within the micrometer range. These alloys were tested under cutting conditions A and B and the results are shown in Table 5.

TABLE 5

Example 12

An alloy powder composed of 2.0% TiC, 6.0% TaC, 5.6% CO and WC residue was formed in shape No.SNG432, sintered at 1450 ° C. in vacuum, and cooled in a similar manner to Example 5. Obtain an alloy with a structure and hardness distribution substantially similar to 5.

Claims (25)

Periodic table IVa, Va, of the periodic table on a cemented carbide substrate composed of at least one hard phase selected from the group consisting of carbides, nitrides and carbonitrides of the Group IVa, Va, and VIa metals and at least one bond phase selected from the iron group metals. In at least one member selected from the group consisting of carbides, nitrides, oxides and borides of Group VIa metals and surface-coated cemented carbide having a single layer or multiple layers of a solid solution thereof, from the interface between the coating layer and the substrate. The hardness of the cemented carbide substrate in the range of ˜5 μm is 700-1300 kg / mm2 with a beaker hardness of 500 g load, and increases uniformly toward the inside of the substrate so as to be constant within the range of approximately 50-100 μm from the interface. Cemented carbide. The coated cemented carbide alloy according to claim 1, wherein the hardness of the cemented carbide substrate within the range of 2 to 5 µm from the interface is beaker hardness of 950 to 1250 kg / mm2. The coated cemented carbide alloy according to claim 1, wherein the hardness of the cemented carbide substrate within the range of 50 to 100 µm from the interface is a beaker hardness of 1500 to 1700 kg / mm 2. The coated cemented carbide alloy according to claim 1, wherein the surface vicinity of the cemented carbide substrate is mainly composed of a layer composed of WC and CO and having a thickness of 5 to 10 µm. The coated cemented carbide alloy according to claim 1, wherein the amount of bonding phase of the cemented carbide substrate in the range of 2 to 20 µm to 50 to 100 µm from the interface is 1.5 to 7 weight times the average amount of the bonding phase present in the alloy. The coated cemented carbide alloy according to claim 1, wherein the amount of bonding phase of the cemented carbide substrate in the range of 2 to 20 µm from the interface is 1.5 to 7 weight times for the range of approximately 50 to 100 µm. The method according to claim 1, wherein the amount of free carbon when the binding phase is made of CO is 1 to 2.4% by weight based on CO, and the amount of free carbon when the binding phase is made of Ni is 0.5 to 2.2% by weight based on Ni. Coated cemented carbide characterized by the above-mentioned. The amount of free carbon [FC] and nitrogen [N] in the cemented carbide substrate is as follows.

(Wherein [FC] and [N] are expressed in weight percent) Coated cemented carbide, characterized in that having. The coated cemented carbide alloy according to claim 1, wherein the amount of the bonding phase in the range from the interface to 5 mu m is smaller than the inside of the cemented carbide substrate. The coated cemented carbide alloy according to claim 1, wherein the content of the bonded phase of the cemented carbide substrate within the range of 3 µm from the interface is smaller than the range below 3 µm from the interface. The cemented carbide substrate according to claim 1, wherein the cemented carbide substrate is composed of 10 to 95% by weight of WC, 1 to 70% by weight of Ti, W and Ta or Nb mixed carbonitride, and 3 to 20% by weight of CO. . The coated cemented carbide according to claim 1, wherein the CO or CO and C in the surface layer are moved by chemical, mechanical or electrochemical treatment. Starting materials that correspond to the components of the hard and bound phases or can form these components in situ are mixed and sintered by decomposition or reaction, and the mixture is cooled at a cooling rate of 0.1 to 10 DEG C / min, A method for producing the coated cemented carbide according to claim 1, comprising coating the resulting cemented carbide substrate with a coating material corresponding to the component. The method of manufacturing a coated cemented carbide according to claim 13, wherein the cooling is performed at a cooling rate of 0.1 to 10 DEG C / min within a temperature range of 1310 to 1225 DEG C. The method of claim 13 wherein the mixture to be cooled has the following relationship:

(Wherein [FC] and [N] are expressed in weight percent) Method for producing a coated cemented carbide, characterized in that it has a content of free carbon [FC] and nitrogen [N] satisfying the. 14. The method for producing a coated cemented carbide alloy according to claim 13, wherein the sintering is carried out while suppressing the denitrification reaction until cooled to 1310 占 폚. 17. The method of claim 16, wherein the denitrification is suppressed by introducing at least one selected from the group consisting of N ₂ , CH ₄ , H ₂ , and Ar. 15. The method of claim 14, wherein the cooling from 1310 DEG C. to 1225 DEG C is carried out in a vacuum or an oxidizing atmosphere. The method of claim 13, wherein the cemented carbide substrate is subjected to chemical, mechanical or electrochemical treatment prior to coating to move CO or CO and C in their surface layer. Starting materials which correspond to the components of the hard and binding phases or which can form these components in situ are mixed and sintered by decomposition or reaction, and the mixture is subjected to a time period of 10 minutes to 15 hours within a temperature range of 1310 to 1225 ° C. A process for producing the coated cemented carbide alloy of claim 1, comprising cooling with a mixture. The mixture of claim 20 wherein the mixture to be cooled has the following relationship:

(Wherein [FC] and [N] are expressed in weight percent) Method for producing a coated cemented carbide, characterized in that it has a content of free carbon [FC] and nitrogen [N] satisfying the. 21. The method of claim 20, wherein the sintering is carried out while suppressing the denitrogenation reaction until cooled to 1310 占 폚. 23. The method of claim 22, wherein the denitrification is suppressed by introducing at least one selected from the group consisting of N ₂ , CH ₄ , H ₂ , and Ar. 15. The method of claim 14, wherein the cooling is carried out in a vacuum or an oxidizing atmosphere. 21. The method of claim 20, wherein the cemented carbide substrate is subjected to chemical, mechanical or electrochemical treatment prior to coating to move CO or CO and C in their surface layer.

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